While the discussions and detailed descriptions of the present invention set forth hereinbelow have been primarily directed toward the deposition of thin films of amorphous semiconductor material, it should be noted that the present technique can be used for the high-rate, low temperature deposition of many other materials. Particularly, the method can be used to deposit materials whose total microstructural local environment includes amorphous, nanocrystalline, microcrystalline, polycrystalline and monocrystalline as well as combinations thereof and intermediate metastable states. The materials which may be deposited by the method of the present invention are not limited to semiconductors, but can also include metals, ceramics, glasses, polymers, dielectrics, insulators and photoactive materials. Of particular interest are semiconductors, high density oxides and nitrides, conductive metals and ceramic protective coatings. Also of interest are nanocrystalline materials having grain features of 30 Angstroms or less.
Another example of a field of particular interest, other than semiconductor materials, is metallic aluminum and/or copper conductive pathway deposition for ultra large-scale integration (ULSI) of electrical circuits. Today, as integrated circuit technology strives for higher and higher integration densities, the width of the electrical current carrying pathway required to achieve these integration densities has become much smaller. In the present state of the technology, there is difficulty with the physical vapor deposition process by which such narrow conducting lines are formed by conventional sputtering techniques. Deposition of conductive pathway material into the openings in the photolithographic masking, which are sometimes as small as nanometers wide, causes a shadowing problem. That is, because PVD produces non-conformal coverage, which leads to thinning at the edges and walls of vias and trenches and to big holes in vias. Therefore, there is a limit to the narrowest effective width to which the conductive pathways can be deposited, without forming open circuits therein.
In general, CVD offers better coverage conformality than PVD. Given the chemical vapor deposition method of the present invention, the depositing species can have a much higher surface mobility, and can therefore "fill in" the shadowed areas, thereby effectively extending the lower limit on the narrowest effective width to which conductive pathways can be deposited.
Prior to turning to the use of the novel method disclosed herein for depositing high quality amorphous silicon alloy material, it is significant to elucidate the reader as to the full capabilities of this method to deposit new and phenomenally different quantum mechanical materials.
A considerable effort has been made to develop processes for readily depositing amorphous semiconductor alloys or films, each of which can encompass relatively large areas, if desired, limited only by the size of the deposition equipment, and which could be readily doped to form p-type and n-type materials where p-n junction devices are to be made therefrom equivalent to those produced by their crystalline counterparts. Amorphous silicon or germanium (Group IV) films are normally four-fold coordinated and were found to have microvoids and dangling bonds and other defects which produce a high density of localized states in the energy gap thereof. The presence of a high density of localized states in the energy gap of amorphous silicon semiconductor films results in a low degree of photoconductivity and short carrier lifetime, making such films unsuitable for photoresponsive applications. Additionally, such films cannot be successfully doped or otherwise modified to shift the Fermi level close to the conduction or valence bands, making them unsuitable for making p-n junctions for solar cell and current control device applications.
A reduction of the localized states was accomplished by glow discharge deposition of amorphous silicon films wherein a gas of silane (SiH.sub.4) was passed through a reaction tube where the gas was decomposed by an r.f. glow discharge and deposited on a substrate at a substrate temperature of about 500.degree.-600.degree. K. (227.degree.-327.degree. C.). The material so deposited on the substrate was an intrinsic amorphous material consisting of silicon and hydrogen. To produce a doped amorphous material a gas of phosphine (PH.sub.3) for n-type conduction or a gas of diborane (B.sub.2 H.sub.6) for p-type conduction were premixed with the silane gas and passed through the glow discharge reaction tube under the same operating conditions. The gaseous concentration of the dopants used was between about 5.times.10.sup.-6 and 10.sup.-2 parts per volume. The material so deposited included supposedly substitutional phosphorus or boron dopant and was shown to be extrinsic and of n or p conduction type.
Amorphous semiconductor alloys made by the prior art processes have demonstrated photoresponsive characteristics ideally suited for photovoltaic applications. These prior art processes however suffered from relatively slow deposition rates and low utilization of the reaction gas feed stock which are important considerations from the standpoint of making photovoltaic devices from these materials on a commercial basis. To alleviate the problem of slow deposition rate, a microwave plasma deposition processes was invented by Ovshinsky, et al. See U.S. Pat. No. 4,517,223, Method of Making Amorphous Semiconductor Alloys and Devices Using Microwave Energy, which issued on May 14, 1985, the disclosure of which is hereby incorporated by reference. The microwave plasma process herein provides substantially increased deposition rates and reaction gas feed stock utilization. Further, the microwave process resulted in the formation of reactive species not previously obtainable in sufficiently large concentrations with other processes. As a result, new amorphous semiconductor alloys could be produced having substantially different material properties than previously obtainable.
The principal advantage that microwave (e.g., 2.45 GHz) generated plasmas have over the more widely used radio frequency (e.g., 13.56 MHz) generated plasmas is that the thin film deposition rates are generally higher by factors of 10 to 100. This enhanced rate comes about because of the increased fraction of all the electrons in the plasma which have the necessary energy to be chemically relevant, i.e., energies greater than about 3 to 4 eV in which to excite molecules, and especially energies greater than 8 to 12 eV in which to dissociate and/or to ionize molecules. In other words, for the same power densities, microwave plasmas are more effective in generating the chemically active species (i.e., excited molecules, dissociated, very reactive radicals, ions, etc.), which are necessary to form thin films, than are radio frequency plasmas. One obstacle that has prevented thin film microwave plasma deposition from being more widely used is that the thin film material properties (e.g., electronic, density, etc.) tend to be inferior to those deposited by radio frequency plasma.
More recently, dual-frequency plasma deposition has been investigated as a method for simultaneously combining the advantageous features of both microwave and radio frequency plasma depositions. That is, while microwave plasma is very efficient for generating active species in the gas phase (because of a relatively higher population of electrons in the energetic tail of the electron energy distribution function), resulting in deposition rates an order of magnitude higher than those observed at lower frequencies, the quality of the deposited material is lacking. On the other hand, radio frequency plasma is characterized by a negative d.c. bias at the cathode, which controls the flux and energy of ions impinging on the cathode surface, resulting in high quality thin film deposition, but at very low deposition rates. Dual-frequency plasma deposition using both microwave energy at 2.45 GHz and radio frequency energy at 13.56 MHz has been shown to produce better quality thin films than microwave energy alone at a higher deposition rate than radio frequency energy alone. See Klemberg-Sapieha et al, "DUAL MICROWAVE-R.F. PLASMA DEPOSITION OF FUNCTIONAL COATINGS", Thin Solid Films, vol. 193/194, December, 1990, pages 965-972.
Although this dual-frequency (DF) plasma deposition method clearly has advantages over single frequency microwave (MW) or radio frequency (RF) plasma, it clearly only a compromise. While DF plasma has higher deposition rates than RF plasma, the rates are unquestionably lower those of MW plasma. Also, while the quality of DF plasma deposited films is greater than those deposited by MW plasma, it is undoubtedly lower than those deposited by RF plasma. It has been found that at the higher deposition rates of MW plasma deposition, it becomes increasingly important that the active neutrals (i.e. chemically activated, electrically neutral species) have higher kinetic energy so that when they arrive at the film surface, they will have the necessary surface mobility to create a high quality film.
In the discussion of microwave plasma deposition above attention should be focused on the electrons, for it is the electrons, which excite, dissociate and ionize the gaseous molecules. One possible mechanism why microwave energy leads to higher electronic energy in the plasma is that the electrons can resonantly couple to the energy. For example, in typical microwave plasmas, the electron density can range from 10.sup.10 to 10.sup.11 per cm.sup.3. At these densities, the resulting plasma oscillation frequency is on the same order as the applied radiation, namely in the low GHz range. Although a tremendous amount of the energy of the high frequency radiation can be pumped into the electrons by resonantly coupling, little if any of that energy can be transferred into kinetic energy of the ions and neutral molecules. This is because the ions are too massive to respond to the high frequency oscillations of the microwave and/or radio frequency energy, and also the energy gain of the electrons cannot be effectively transferred to the ions/neutrals due to the their large difference in mass.
There are two ways to resonantly couple energy into a plasma: (a) by high frequency radiation, which couples into the electrons; and (b) by low frequency radiation, which couples into the ions. When energy is pumped into the ions, their energy is very effectively transferred into the neutrals owing to their similarity in masses. Heating of the neutrals is desirable because it is the neutrals which overwhelmingly account for the deposition rate. It should be noted that because a relatively few ions must heat a large number of neutrals, efficiency requirements dictate the need to resonantly couple energy into the ions.
Large scale uniform deposition in rf plasma systems is not a serious issue because the wavelength at 13.56 MHz is more than 22 m, i.e., much greater than any practical deposition chamber. However, since the wavelength at the microwave frequency of 2.45 GHz is only 12.2 cm, the technical challenges in achieving uniform depositions over the desired scale of 30 cm are evident. Because of this, deposition results (e.g., film quality from a certain point in parameter space) obtained from small sized microwave research reactors (.about.5 cm substrates), cannot be automatically applied to the larger production systems because the hardware implementation will be fundamentally different. Scaling the process to larger sizes is not simply a matter of making larger electrodes (as it is in rf plasma systems) but involves new concepts. Although others have demonstrated high quality a-Si:H films with solar cell efficiencies approaching the best rf deposited material, similar results on the large production scale have yet to be demonstrated.
The record efficiency of 13.7% (active area, initial efficiency) for a multijunction a-Si alloy cell from PECVD was reported in 1988 by J. Yang et al. of ECD Proceedings of the 20th IEEE PV Specialists Conference (1988) page 241. Since then, progress has only been made in improving the large area module efficiency, which has now reached the initial efficiency of 10-11% and stable efficiency of 9-10%. The lack of progress in improving the quality of the narrow bandgap (E.sub.g &lt;1.5 eV) a-SiGe alloy as well as the wider bandgap a-SiC alloys (eV&gt;1.8 eV) is the primary reason preventing further efficiency enhancement. In order to reach the DOE goal of a stable 15% efficiency for modules by the year 2005, a major breakthrough of materials on either narrow bandgap a-SiGe alloy material or wider bandgap a-SiC alloy materials has to be made. The glow discharge technique which has been very successful in fabricating high quality a-Si:H has not been able to provide equivalently high quality material where the base alloy has been modified to widen or narrow the bandgap. Simply stated, the photovoltaic properties of the alloys severely deteriorate with increasing Ge or C content.
Generally, the film quality of a-SiGe alloys, as characterized by the defect density measured from sub-bandgap absorption, the network disorder measured from the Urbach parameter and the photosensitivity (.sigma..sub.light .sigma..sub.dark), deteriorate with increasing Ge content. The deterioration of the material quality severely affect the fill factor (FF) of the devices. As the Ge content of the intrinsic layer of a PV device exceeds 40% (i.e., E.sub.g &lt;1.5 eV), the FF becomes so poor that no decent device can be made. To maintain a good FF (.about.0.7), the current multijunction solar cell can only use materials with 1.5 eV.ltoreq.E.sub.g .ltoreq.1.75 eV. Should the narrow bandgap 1.2 eV&lt;Eg&lt;1.5 eV a-SiGe alloy be made of a quality which approaches that of an a-Si:H alloy, a multijunction laboratory device would have an efficiency of .gtoreq.17% (i.e. Voc=2.5 V, FF=0.75 and Jsc&gt;9 mA/cm.sup.2) and a 15% stable module efficiency could be obtained.
In a conventional PECVD system, the gas phase environment contains electrons, various types of ions, chemically active neutral species (i.e. free radicals) and molecules. At an operating pressure of 0.5-1.0 Torr the degree of ionization is extremely low (i.e. &lt;10.sup.-6). Therefore, only the free radicals contribute to the film growth. Electrons, due to their light mass, gain energy from the electric field much more effectively than the heavier ions. The electron energy distribution can be roughly described as a Maxwell-Boltzmann distribution with average electron energies around 2 eV, corresponding to an electron temperature of .about.10.sup.4 K. Since the ions are able to receive some energy from the field, their temperature is slightly above ambient. The typical ion energy is around 0.04 eV or .about.500 K. Because of the low degree of ionization, and the low ion energy, the neutral free radicals remain essentially at ambient temperature. In other words, the kinetic temperature of the film precursors from conventional PECVD processes is around 0.025 eV or 300 K.
In any deposition process, the surface mobility of the precursor species is a critical parameter in determining the quality of the depositing film. It is important that the precursor species, upon their arrival at the surface, be able to move around to find energetically favored sites and to fill microvoids. Since the film precursors from PECVD are not energetic, the precursor primarily rely on the substrate temperature (T.sub.s) as the energy source. Unfortunately, the T.sub.s of hydrogenated amorphous films are limited by the hydrogen evolution temperature which is much lower than the crystallization temperature. For example, the crystallization temperature of silicon is .about.630.degree. C., whereas hydrogen evolution from the a-Si:H film begins at a temperature of about 250.degree.-275.degree. C. Films deposited at T.sub.s above 275.degree. C. from a conventional PECVD system will have a higher number of dangling bonds due to hydrogen evolution. In the case of a-SiGe:H film deposition, because the bond strength of Ge-H is weaker than Si-H, hydrogen evolution from Ge begins at 150.degree. C. To prevent the preferential loss of hydrogen, one would expect that the film should be deposited at a T.sub.s of .about.150.degree. C. However, experiments indicate that device performance of a-SiGe film deposited at T.sub.s of .about.150.degree. C. is rather poor. Most of the so called "device quality" a-SiGe alloys with E.sub.g .gtoreq.1.5 eV are actually deposited at .about.300.degree.-325.degree. C. As a result, the narrow bandgap films have a higher dangling bond density than the a-Si:H films.
Why do the a-SiGe alloys prefer to grow at higher temperatures, despite hydrogen evolution? One possibility is that Ge precursors in the plasma are less mobile than Si precursors because the Ge atom is bigger and heavier than the Si atom. In order for Ge to move as freely as Si on the growing surface, a higher T.sub.s is desirable. If not for H evolution, higher quality a-SiGe films might be grown at even higher temperatures. This argument also explains why the quality of the depositing film deteriorates with increasing Ge content, i.e., because at temperatures of 300.degree.-325.degree. C., the preferential loss of H from Ge sites is aggravated.
The lack of surface mobility in the depositing a-Ge:H film is also manifested by the high degree of disorder in the network and high porosity present in the deposited films. If the precursor species were mobile, they would have a greater opportunity of finding low energy sites so that not only a denser film could be made, but better, stronger bonds could be formed. Thus, increasing surface mobility of depositing species will lead to a reduction of weak bond formation and elimination of the Stablaer-Wronski effect. In the case of a-Si:H film growth, high quality material is typically deposited at a rate of 1-2 .ANG./sec at 250.degree. C. The quality of the film deteriorates with increasing deposition rate. Since surface diffusion is a rate-limiting process, this result indicates that at 250.degree. C., even the Si precursors can barely move fast enough to accommodate even low rate deposition. Should the surface mobility of Si precursors be further enhanced without increasing T.sub.s, it would enable high quality a-Si:H alloy material to be deposited at much higher rates, even up to 100 .ANG./sec. High rate deposition translates into high production throughput and lower production costs.
From the above discussion, it becomes clear that one of the fundamental problems in the deposition of high quality a-SiGe films is fighting the conflicting growth parameters of surface mobility and hydrogen evolution.
Since fluorine bonds to Ge much more strongly than H, one would expect that the temperature fluorine evolution from Ge should be much higher than that of H from Ge. This is one reason to use a fluorine containing gas as precursor in the a-SiGe alloy deposition. However, a fluorine passivated surface is not as "slippery" as a H covered surface, because F is so reactive that it bonds too readily with film precursors at the first point of contact. Therefore, the a-SiGe:F film has more network disorder than an a-SiGe:H film. Nevertheless, it has been demonstrated that small amounts of F does help in making higher quality films.
Since there has not as yet been a way found to use F to replace H, and since the substrate temperature cannot be raised without H evolution, the only viable solution for depositing high quality a-SiGe film seems to be raising the kinetic temperature of the precursor species. As was pointed out previously, the kinetic temperature of the precursor species in PECVD is virtually ambient and the species will be heated at most to the substrate temperature. Clearly this temperature is not high enough.
Described hereinbelow are several unique features of PECVD which are important to a-Si and a-SiGe growth. First, there is an abundance of atomic hydrogen in the plasma. The atomic hydrogen on the growing surface can remove weak bonds and passivate dangling bonds. Hydrogen is needed in the structure not only to passivate dangling bonds, but to release strain in the tetrahedrally bonded amorphous network. Secondly, the dominant film precursor for a-Si:H deposition from PECVD is generally believed to be SiH.sub.3. This is due to the low reactivity of SiH.sub.3 with the primary discharge gases such as H.sub.2 and SiH.sub.4, which allows it to diffuse to the substrate surface. The advantage of such a low reactivity precursor is that it will have a small probability of surface reaction at the initial point of contact, and thus requires less energy to move around on that surface. This may be one of the reasons that PECVD a-Si:H exhibits the best photovoltaic properties compared to films deposited by other techniques. By analogy to silane chemistry, it is expected that GeH.sub.3 is the dominant precursor for a-Ge:H film. However, GeH.sub.3 is much heavier and diffuses slower than SiH.sub.3. Hence, an a-Ge:H film is much more porous and of poorer quality than an a-Si:H film under the same deposition conditions. Finally, even though free radicals such as SiH, SiH.sub.2, SiH.sub.3, etc. are unable to gain kinetic energy, through elastic collisions with electrons, they can acquire quite high internal energy, in the form of vibrational and rotational energy, through various inelastic collision processes. For example, the vibrational and rotational temperatures of the free radical SiH have been measured from the intensity of the optical emission spectrum. In the electronically excited state, the temperatures are 4000 K. and 1800 K., respectively. In the ground state, they are 2000 K. and 485 K., respectively. The high internal energy of free radicals is another special feature of PECVD. It had been speculated that because of the high internal energy of free radicals, microcrystalline films can be grown at temperatures below 200.degree. C.
In a-Si:H and/or a-SiGe:H film growth, the precursor species such as GeH.sub.3 and/or SiH.sub.3 arrive at the surface, move around, and locate the most energetically favorable low energy sites, thus forming a thin film on the substrate. Hydrogen elimination and a restructuring of the surface bonds will proceed to complete the growth process. However, energy is required for these processes to occur. The high internal energy of the precursor species may be an important energy source for the successful initiation of these processes. Therefore, techniques which solely heat the gas to increase the translational temperature of the precursor species cannot ensure the fabrication of high quality film. A number of aspects which are critically important to film growth in PECVD can be identified. For example, the abundance of atomic hydrogen, the low reactivity of precursor species and the high internal energy of those species are all important attributes in the growth of a-Si:H and a-SiGe:H films and should be present in the any advanced PECVD technology. A drawback with conventional embodiments of PECVD is the low kinetic temperature of the precursor species which have prevented the growth of high quality a-SiGe alloy films.
Therefore, there is still a strong felt need in the art for a safe method for the high-rate, low-temperature deposition of high-quality materials on large area substrates.